Case hardened dies for improved die life

ABSTRACT

Steel alloys susceptible to case and core hardening comprise 0.05 to 0.24 weight percent carbon; 15 to 28 weight percent cobalt and 1.5 to 9.5 weight percent in nickel, small percentages of one or more additives: chromium, molybdenum, and vanadium; and the balance iron. Carburized roll form and punch dies and tools made from case hardened steel alloys with a reduced hardness core provide high wear and fatigue resistance as well as improved contact and bending fatigue resistance thereby avoiding premature failure and extending the useful life of such tools.

CROSS REFERENCE TO RELATED APPLICATION

This is a continuation in part of application Ser. No. 09/664,021 filedSep. 19, 2000 entitled “Case Hardened Dies For Improved Die Life” whichis a continuation in part of application Ser. No. 09/239,131 filed Jan.28, 1999, now U.S. Pat. No. 6,176,946 entitled “Advanced CaseCarburizing Secondary Hardening Steels” which is a continuation ofProvisional Application Serial No. 60/072,834 filed Jan. 28, 1998entitled “Advanced Case Carburizing Secondary Hardening Steels”, and forwhich priority is claimed with respect to said applications.

ACKNOWLEDGMENT OF FUNDING

The original subject matter of the parent application was funded, atleast in part, by the Army Research Office, Grant No. DAAH04-96-1-0266.The subsequent continuation-in-part subject matter herein was not fundedby Grant No. DAAH04-96- 1-0266.

BACKGROUND OF THE INVENTION

This invention relates to a new class of steel alloys especially usefulfor the manufacture of case hardened steel gears and other products, inparticular, blades such as skating blades made from case hardened steelalloys.

Currently, there are a number of high performance steels on the market.Many of these materials utilize primary carbides to achieve high surfacehardness and others use stage one or stage three tempered conditionswith epsilon carbide or cementite strengthening. Primary carbides areformed when the carbon content exceeds the solubility limit during thesolution treatment, and large alloy carbides precipitate. This is thecase for secondary hardening steels using alloy carbide strengtheningfor greater thermal stability to improve properties such as scoringresistance.

However, primary carbide formation can have a detrimental impact on bothbending and contact fatigue resistance. Formation of primary carbidescan also make process control difficult for avoidance of undesirablecarbide distributions such as networks. In addition, primary carbideformation in current steels can lead to a reversal in the beneficialresidual compressive stresses at the surface. This is due to a reversalof the spatial distribution of the martensite start temperature due tothe consumption of austenite stabilizing elements by the primarycarbides. In applications of sliding wear, the formation of primarycarbides may be beneficial.

Thus, there has developed a need for case hardenable steel alloys whichdo not rely upon primary carbide formation, but which provide secondaryhardening behavior for superior thermal stability. This inventionprovides a new class of steel alloys meeting this requirement, whileexploiting more efficient secondary hardening behavior to allow highersurface hardness levels for even greater improvements in fatigue andwear resistance. Rolls for manufacturing processes utilizing such steelsare projected to have more uniform and enhanced performancecharacteristics derived through simplified manufacturing technologiesand to also have performance characteristics which are more predictableand reproducible and lead to longer life of the tool.

SUMMARY OF THE INVENTION

Briefly, the present invention comprises a class of case hardenablesteel alloys in the form of roll form dies with carbon content in therange of about 0.05 weight percent to about 0.24 weight percent incombination with a mixture of about 15 to 28 weight percent cobalt, 1.5to 9.5 weight percent nickel, 3.5 to 9.0 weight percent chromium, up to3.5 weight percent molybdenum, and up to 0.2 weight percent vanadium.

The microstructural features are a Ni—Co lath martensite matrix steelstrengthened by M₂C carbides typically containing Cr, Mo and V. Typicalprocessing of this class of steels includes case carburizing, solutiontreatment, quenching, and tempering, although due to the high alloycontent, quenching may not be required. Case carburizing produces agradient in the volume fraction of the M₂C carbides and results in aconcomitant increase in hardness and promotes a surface residualcompressive stress. The efficiency of the M₂C strengthening responseallows this class of steels to achieve very high surface hardnesses withlimited soluble carbon content. Thus, this class of steels have theability to achieve very high surface hardnesses without the formation ofprimary carbides.

Typical advantages of this class of alloys include ultrahigh casehardness leading to superior wear and fatigue resistance, superior corestrength and toughness properties, optional air hardening resulting inless distortion, and higher thermal resistance.

This new class of secondary hardening steels are matrix steels utilizingan efficient M₂C precipitate strengthening dispersion. Because of theefficiency of this strengthening dispersion, a superior combination ofproperties can be attained for many applications on a situation bysituation and product by product basis. For example, in situations wherethe desired surface properties are similar to current materials, thecore strength and toughness can be superior. In applications wheresuperior surface properties are desired, the disclosed steels can easilyoutperform typical materials while maintaining normal core properties,and in applications which require corrosion resistance, these new steelscan provide stainless properties with surface mechanical propertiessimilar to typical non-stainless grades.

Thus, an object of the invention is to incorporate desirable propertiesresulting from the class of alloys disclosed in various products.

A further object of the invention is to provide roll form dies, diepunches, stamping dies, or draw dies made from case hardened steel alloymaterials wherein the surface hardness of the dies surface and the corehardness of the dies are controlled to maximize performance and toprovide uniform and reproducible characteristics.

Another object is to provide die forms for rolling, punching, stampingor drawing made from steel alloy materials disclosed wherein theflexibility, hardness, sharpness of the die and other diecharacteristics are controllable and reproducible.

These and other objects, advantages and features of the invention willbe set forth in the detailed description which follows.

BRIEF DESCRIPTION OF THE DRAWING

In the detailed description which follows, reference will be made to thedrawing comprised of the following figures:

FIG. 1 is a graph correlating hardness to precipitation driving forcefor experimental and predicted results;

FIG. 2 is a graph correlating precipitation half completion time andhalf completion coarsening rate constant for experimental and predictedresults;

FIG. 3 is a graph which correlates calculated segregation free energydifference with experimental embrittlement potency;

FIG. 4 is a flow block diagram of the total system structure of thealloys of the invention;

FIG. 5 is a graph depicting the relationship between cobalt and nickelcontent for a 200° C. Ms temperature for the alloys of the invention;

FIG. 6 is a pseudo-ternary diagram as a function of chromium, molybdenumand vanadium at 0.55 weight percent carbon with regard to alloys of theinvention at 1000° C.; and

FIG. 7 is a graph comparing hardness of the steel alloys of theinvention with conventional carburized alloys.

FIG. 8 is a graph containing Falex wear test data for steel alloys ofthe invention in comparison with conventional 8620 steel.

FIG. 9 is a graph of NTN 3 ball-on-rod rolling contact fatigue data foralloys of the invention in comparison with conventional M50 bearingsteel;

FIG. 10 is an isometric view of a typical roll form die fabricated froma carburized case hardened steel alloy; and

FIG. 11 is a cross-sectional view of the die of FIG. 10 taken along theline 11—11.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

The steel alloys of the invention were developed using various modelingtechniques followed by experimental confirmation or testing. Animportant component of the modeling is the application of athermochemical data bank and software system. The system or programemployed uses thermodynamic assessments from binary, ternary, andquaternary systems to extrapolate to higher order multicomponentsystems. Equilibria, constrained equilibria, and driving forces can becalculated as functions of composition, chemical potential, as well asother user defined functions. To apply this information to the modelingof highly nonequilibrium processes of interest in real alloys, thedynamic nature of phase transformations in terms of thermodynamicscaling factors are described and then evaluated by the thermochemicalsoftware. Thus, hypothetical steel compositions were the subject of aninitial computational model involving the precipitation of M₂C carbidesleading to a secondary hardening response in ultrahigh-strength steels.A second effort employed a published thermodynamics-based model for thenon-linear composition dependence of the martensite start temperature. Athird modeling effort involves the application of quantum mechanicalcalculations to the production of hypothetical compositions with thegoal of achieving improved resistance to hydrogen embrittlement andintergranular fracture. Modeling techniques were then followed bytesting of the optimized alloys. Following is a discussion of modelingtechnique considerations.

Secondary Hardening

Ultrahigh-strength (UHS) secondary hardening steels are strengthened bythe precipitation of coherent M₂C carbides during tempering. In high Costeels in which dislocation recovery is retarded, the M₂C carbidesprecipitate coherently on dislocations and provide the characteristicsecondary hardening peak during tempering. A wide range of techniquesare utilized to gather experimental information across a complete rangeof size and time scales of interest. Atom-probe field-ion microscopy(APFIM), transmission electron microscopy (TEM), small angle neutronscattering (SANS), and X-ray diffraction (XRD) techniques provideinformation on the size, shape, composition, and lattice parameters ofthe M₂C precipitates as well as the resulting hardness values spanningtempering times of less than an hour to more than a thousand hours. Thisstudy identified that the precipitation was well described by a theorydeveloped by Langer and Schwartz for precipitation at highsupersaturation in which the growth regime is suppressed andprecipitation occurs by a process of nucleation and coarsening,maintaining a particle size close to the critical size.

Based on these investigations, two important scaling factors areidentified. The initial critical nucleus determines the size scale ofthe precipitates throughout the precipitation reaction and thecoarsening rate constant determines the precipitation time scale. Thepeak hardness in an ultra high strength steel commonly occurs at theparticle size corresponding to the transition from particle shearing toOrowan bypass. It is also advantageous to bring the M₂C precipitation tocompletion in order to dissolve all of the transient cementite whichotherwise limits toughness and fatigue life. Therefore, the smaller theinitial critical particle size, the closer completion of precipitationoccurs to peak hardness and more efficient strengthening is obtained.The time scale of precipitation is also important due to the kineticcompetition between the secondary hardening reaction and the segregationof impurities to the prior austenite grain boundaries leading tointergranular embrittlement.

The initial critical nucleus size scales inversely with thethermodynamic driving force for precipitation. In the case of the M₂Ccarbide it is important to include the influence of prior cementiteformation and coherency on this quantity. The coherency elastic selfenergy can be evaluated by the calculation of an anisotropic ellipsoidalinclusion using the equivalent Eigenstrain method and the impact ofsolute redistribution on the resulting stress distribution is addressedby using open-system elastic constants. By relating the coherency strainto composition via the compositional dependence of the particle andmatrix lattice parameters, the composition dependence of the elasticself energy is determined in a form compatible with the thermodynamicsoftware. The linear elastic self energy calculation represents an upperlimit and a correction factor is used to fit the precipitationcomposition trajectories of a large set of experimental alloys.

The impact of prior cementite precipitation is accounted for by thecalculation of the coherent driving force in the presence of the carbonpotential due to para-equilibrium cementite. This paraequilibrium carbonpotential is defined by an equilibrium between the matrix and cementitein which the substitutional species are held constant and only theinterstitial carbon is allowed to partition. In this approximation thecementite acts as a carbon source at constant chemical potential.

FIG. 1 represents the level of agreement of the strengthening responseof the model alloys with the above model. The model alloys contain 16 wt% Co, 5 wt % Ni, and 0.24 wt % C with varying amounts of the carbideformers Cr, Mo, and, in a few cases, V. The nickel content is chosen toeliminate austenite precipitation during tempering which otherwisecomplicates the hardening response. In FIG. 1, the peak hardness duringtempering at 510C is plotted against the driving force for precipitationof the coherent M₂C carbide in the presence of para-equilibriumcementite. The open circles represent alloys containing V. Therelationship demonstrates the ability to predict peak hardness valueswithin approximately +/−25 VHN in this class of alloys.

The time scale of precipitation at high supersaturations, according tothe Langer-Schwartz treatment, scales with the coarsening rate of theparticle distribution. The modeling pursued in this work expands uponthe Lifshitz-Slyozov and Wager (LSW) theory, describing the coarseningof spherical particles in a binary system, with the intent of removingthe binary restrictions of the LSW theory and reformulating it in amanner compatible with the multicomponent thermodynamics of the softwareand data bank system.

The result of this analysis characterizes the coarsening rate of aparticle of average size as a function of the multicomponent diffusioncoefficients, the equilibrium partitioning coefficients, and the secondderivatives of the Gibbs free-energy evaluated at the equilibrium state.The surface energy and molar volume are taken to be compositionindependent and are considered constant. In this form, the coarseningrate constant is the result of an asymptotic analysis and is onlyrepresentative at very long time scales and very close to equilibrium.This is certainly not the case for the precipitation of the M₂C carbideat high supersaturation. The matrix content of alloy is far fromequilibrium during much of the precipitation process, approachingequilibrium only near completion. This effect is more severe for alloyscontaining stoichiometric quantities of carbide formers as measured bythe relative difference in the matrix alloy content during precipitationand at equilibrium.

During precipitation in a stoichiometric alloy, the alloy matrix contentis of the same order as the overall alloy content, while at equilibrium,the matrix alloy content is very small. To define a coarsening rateconstant more representative of the conditions present during theprecipitation process, a coarsening rate is evaluated at the point whenthe volume fraction of the precipitate is one-half of the equilibriumvalue. This is achieved by calculating the coherent equilibrium for theM₂C carbide, and then, adding energy to the M₂C phase to account forcapillarity until the amount of the phase is half of the equilibriumvalue. The coarsening rate is then calculated from the thermodynamicproperties of this state. FIG. 2 represents the correlation between theprecipitation half-completion time and the half-completion coarseningrate constant of the model alloys for which this data is available.

Ms Temperatures

To predictively control the spatial distribution of martensite-start(Ms) temperatures in the carburized steels to achieve fully martensiticstructures with controlled residual stress distributions, a publishedmodel was employed. The thermodynamics-based nucleation-kinetic modelwas calibrated to the composition-dependence of measured Ms temperaturesusing both literature data and assessments of experimentalmulticomponent alloys.

Interfacial Cohesion

Intergranular embrittlement phenomena such as hydrogen embrittlement areundesirable in the intended alloys. Embrittlement of ultrahigh-strengthsteels is associated with the prior segregation to the grain boundary ofimpurities such as P and S. A thermodynamic treatment of this phenomenonby Rice and Wang illustrates that the potency of a segregating solute inreducing the work required for brittle fracture along a boundary islinearly related to the difference in the segregation energy for thesolute at the boundary and at the free surface. Specifically, a solutewith a higher segregation energy at the free surface will be anembrittler while a solute with a higher segregation energy at the grainboundary will enhance intergranular cohesion. A survey of reportedsegregation energies and embrittling potency (reported as the shift inthe ductile-to-brittle transition temperature per atomic percent soluteon the grain boundary) in Fe-base alloys demonstrates these generaltrends; however, the experimental difficulty of surface thermodynamicmeasurements gives ambiguous values for some solutes.

First principle calculations were used to determine the total energy ofatomic cells representing the Fe υ3 [110](111) grain boundary and (111)free surface with a monolayer of an impurity solute present. Thecalculations were accomplished with the full-potential linearized planewave (FLAPW) total energy technique. The atomic structure in each casewas relaxed to find the minimum energy state. The results of thesecalculations include not only the segregation energies responsible forthe embrittling or cohesion enhancing effects of segregating solutes,but the underlying electronic structure of the solutes in the boundaryand surface environments. A comparison of the directional covalentelectronic structure between B, a strong cohesion enhancer, and P, astrong embrittler, indicates the strong bonding of the B atom across theboundary plane associated with hybridization of the B 2p electrons withthe Fe d band. This directional bonding is not seen in the case of the Patom which does not significantly hybridize with Fe.

The results of the first principle calculations were correlated to theexperimental embrittling potency in FIG. 3. The difference between grainboundary and free surface segregation energies, calculated by electronicstructure calculations, and the experimentally observed shift in theductile-to-brittle-transition-temperature are plotted for C, B, P and Ssolutes. The C and B are shown as cohesion enhancers, P and S asembrittlers. The computed energy differences are in excellent agreementwith the observed effects on interfacial cohesion.

MATERIALS DESIGN

Background

Design considerations for high performance rolls, punches, stamping anddraw dies for various applications include the desire to provide veryhigh surface hardness while maintaining material ductility for shock andflow tolerance. Critically, the portion of the tool (die) which contactsthe work surface must be very hard, i.e., greater than 58 Rockwell Chardness. Heretofore, monolithic tool steels with a hardness rangingfrom about 58 to 62 Rockwell C hardness have been used. However, suchsteels are brittle and do not exhibit a long term work life span.Alternatively, a reduced hardness alloy may be coated with wearresistant or hard coatings to improve tool life and surface lubricity.This involves potentially expensive coating techniques. Alternatively,very high hardness ceramic or powder metal dies can be used, but theyare very expensive and have limited ductility making them hard tomanufacture and prone to damage if not handled properly. The alternativeof the present invention provides benefits not available from theseprior techniques. It involves case hardening by carburizing or nitridingan alloy taken from a special group of alloys in the form of the rollform die, punch, stamping or draw die wherein the form is hardenedgreater than about 58 Rockwell C with the core maintained at a lesserhardness and wherein the surface is hardened at a depth of at leastabout I millimeter in the region of the die wear surface. The core istypically less than about 53 Rockwell C.

Punch dies or stamping dies may also be made in accord with thedescribed procedure and formulation. The work surface will again be casehardened by a carburizing, or nitriding process, or both, typically ofthe select steel disclosed thereby providing a gradient between a 1 to 3millimeter layer of high hardness (>62 Rockwell C) work surface and arelatively softer core or tool matrix (typically <53 Rockwell C).

Analysis

The systems analysis of the case-core secondary hardening steel systemis the first step in the design process. FIG. 4 illustrates the totalprocessing/structure/properties/performance system structure for highpower-density gears manufactured by three alternative processing routes,conventional forged ingot processing, near net shape casting and powderprocessing. Case hardenable secondary hardening gear steels are asubsystem of this flow-block diagram and are the focus of thisdisclosure. The sequential processing steps dictate the evolution of thecase and core microstructures which determine the combination ofproperties required for the overall performance of the system. Both thecase and core consist of microstructures of a martensite with high Co,for dislocation recovery resistance essential for efficient secondaryhardening, and Ni, for cleavage resistance. Strengthening is provided bythe coherent precipitation of fine M₂C carbides on dislocations. Thissecondary hardening reaction dissolves the transient cementite and itbrings the precipitation reaction to completion in order to eliminatecementite for high toughness and fatigue resistance. The grain refiningdispersion has a double impact on toughness. By limiting grain growth athigh temperature during solution treatment, brittle intergranular modesof fracture are inhibited.

The grain refining particles also play an important role in the ductilemicrovoid nucleation and coalescence fracture behavior. Thus, it isdesired to preserve adequate volume fraction and size to pin the grainboundaries while choosing the phases with greatest interfacial cohesion.Also desirable is the control of the grain boundary chemistry to avoidintergranular embrittlement (such as by hydrogen embrittlement) inassociation with prior segregation of embrittling impurities. Duringtempering, impurities segregate to the grain boundaries and in the caseof P and S reduce the interfacial cohesion of the boundary promotingintergranular embrittlement. A number of methods are used to avoid thisproblem. Gettering compounds can be utilized to tie up the impurities instable compounds reducing the segregation to the grain boundary. Inorder to produce the most stable compounds, however, rapidsolidification processing is required. Within their solubility limits,additional segregants such as B can be deliberately added to enhanceintergranular cohesion, and the precipitation rate for the secondaryhardening reaction can be increased to limit the time at temperature forharmful grain boundary segregation.

Design

As a first design step, core and case carbon levels required for thedesired hardness are estimated. This is done by fitting data forexisting secondary hardening Ni—Co steels to an Orowan strengtheningmodel and extrapolating to the desired strength. It is estimated that acore carbon content of 0.25 wt % and a case carbon content of 0.55 wt %is needed to provide the desired core and case hardness in this Ni—Costeel.

The next step is to determine the matrix composition of the Fe, Ni, andCo. In order to produce the desired lath martensite morphology, an M_(s)temperature of 200° C. or above is required. Using the nucleationkinetic model for the compositional dependence of the M_(s) temperature,the variation with Ni and Co content is determined. This result isillustrated in FIG. 5 for the case carbon content using a preliminarycomposition of carbon formers equal to 5 wt % Cr, 0.5 wt % Mo and 0.0 wt% V. Since the case has a higher carbon content than the core, the corewill possess a higher M_(s) temperature than the case. In FIG. 5 the Coand Ni content required to fix the_(s) M temperature at 200° C. isindicated. Since a high Ni content is desired to avoid cleavagefracture, a composition containing 25 wt % Co was chosen. This allowsthe highest possible Ni content, approximately 3.5 wt %, to be used.These calculations are later repeated for consistency when thecomposition of the carbide formers is further refined.

To define the optimal composition of the carbide formers a number ofdesign constraints are applied. The total amount of carbide formers inthe alloy must be greater than that required to consume the carbonpresent in the case. This lower limit insures that, at completion,embrittling cementite is completely converted to M₂C carbide. In orderto reduce grain boundary segregation, the precipitation rate ismaximized. This allows the shortest possible tempering time. Thecoherent precipitation driving force is maximized to provide a smallcritical particle size for the M₂C and more efficient strengthening.Finally, the solution temperature is limited to 1 000° C. This allowsCr, Mo and V containing carbides such as M₂₃C₆, M₇C₃, MC and M₆C to bedissolved at reasonable processing temperatures while maintaining veryfine scale TiC carbides to act as the grain refining dispersion.

Calculations for the precipitation rate constant indicate low Mocompositions are favorable, while driving force calculations havedemonstrated the highly beneficial effect of higher V contents. Thesolubility constraints are presented by the diagram in FIG. 6. Here theequilibrium phase fields at 1000° C. are given as a function of Cr and Vcontent. The Mo content is determined by the stoichiometry requirements,the matrix composition is taken from the earlier calculations, and thecarbon content represents the case composition. The point on the diagramwithin the single phase FCC field that maximizes the V content andminimizes the Mo content represents the compromise fulfilling the designcriteria. This composition is 4.8 wt % Cr, 0.03 wt % Mo, and 0.06 wt %V. A recalculation of the matrix composition using the final carbideformers results in an alloy composition of Fe −25 Co−3.8 Ni−4.8 Cr−0.03Mo−0.06 V −0.55 (case)/0.25 (core) C (in wt %). Consistent with themodel predictions of FIG. 3, a soluble boron addition of 15-20 ppm isadded to enhance intergranular embrittlement resistance.

EXAMPLES

A 17 lb. vacuum induction heat of the above composition was preparedfrom high purity materials. The ingot was forged at 1150° C. in a bar1.25″ square by 38″ long. The M_(s) temperature of the alloy wasdetermined from dilatometery and found to agree with model predictions.The solution treatment response of the alloy was determined fromhardness measurements in the stage I tempered condition. The optimumprocessing conditions for the core material was determined to be a 1050°C. 1 hour solution treatment followed by an oil quench and a liquidnitrogen deep freeze. After optimal solution treatment, a 12 hour temperat 482° C. results in the desired overaged hardness of 55 R_(c) for thecore material. The material was then plasma carburized and processedusing these parameters. The C potential, temperature and time used inthe carburizing treatment were determined from simulations withmulticomponent diffusion software to provide the target surface carboncontent of 0.55 wt % and a I mm case depth. The curve labeled C2 in FIG.7 represents the hardness profile achieved for the carburized sample. Asurface hardness of 67 HR_(c) and a case depth of 1 mm are obtained.

Using techniques and processes of this nature, the following alloys setforth in Table 1 were developed and tested:

TABLE 1 Alloy Fe Co Ni Cr Mo V C (Core) A1 Bal. 18 9.5 3.5 1.1 0.08 0.20C2 Bal. 25 3.8 4.8 0.03 0.06 0.237 C3 Bal. 28 3.25-3.15 5.0 1.75-2.500.025 0.05-0.18 CS1 Bal. 15 1.5 9.0 0.0 0.2 0.05-0.20

The first, A1, is targeted as a replacement for current gear materialsin applications where component redesign is not feasible but higher corestrength and toughness is needed. As such, A1 has surface wearproperties similar to current commercial properties, but possessessuperior core toughness and strength 54 HRC and a K_(IC) of >75 Ksiin.The second alloy C2 corresponds to the prototype alloy just described.The third alloy, C3, pushes the surface properties s far as possiblewhile maintaining adequate core strength and toughness. As also shown bythe hardness profiles of FIG. 7, the alloy has reached a surfacehardness corresponding to HRC 69. Wear tests for the carburized materialin a standard Falex gear simulator show much reduced weight losscompared to standard carburized 8620 steel in FIG. 8. A ball-on-rodrolling contact fatigue test (NTN type) conducted at 786 ksi Hertziancontact stress indicates an order of magnitude increase in L₁₀ fatiguelife compared to M50 bearing steel as shown in FIG. 9. The fourth alloyin Table 1, CS1, represents a stainless variant of this class of alloy.Targeted to match the surface properties of standard non-stainless gearand bearing materials with sufficient core strength and toughness, thealloy has achieved corrosion resistance better than 440° C. by anodicpolarization conducted in distilled water with a neutral ph. (sucroseadded for electrical conductivity). Similar relative behaviors weredemonstrated in 3.5% NaCl solution. In salt fog tests, CS1 outperformed440C and commercial carburizing stainless steels, the performance gapwidened when the tests were completed on samples in the carburizedcondition. The carburized alloy achieved surface mechanical propertiesequivalent to A1 while maintaining corrosion resistance. In RFC tests ofthe type represented in FIG. 9, both A1 and CS 1 showed L₁₀ fatigue lifeequal or superior to the M50 bearing steel. The table (Table 2) belowsummarizes the performance achieved in the four alloys:

TABLE 2 Rolling Core Core Contact Corrosion Hard- Tough- Surface BendingFatigue⁵ Resist- Alloy ness¹ ness² Hardness³ Fatigue⁴ L₁₀ ance⁶ A1  54R_(c) >75 >61 R_(c) ≧EN- ≧M50 NA ksi✓in 36C C2 ≦58 R_(c) adjustable  67R_(c) NA in process NA C3 ≦59 R_(c) adjustable  69 R_(c) NA 10 × M50 NACS1 ≦53 R_(c) ≧25  63 R_(c) NA ≧M50 >440C ksi✓in adjustable ¹Hardnessdetermined by ASTM E18. ²Core toughness determined by ASTM E813.³Surface hardness determined by E384. ⁴Bending fatigue determined using4-point bend testing. ⁵Rolling contact fatigue determined by NTN 3ball-on-rod techniques. ⁶Corrosion resistance determined by anodicpolizations and salt for testing.

Each alloy has a surface hardness exceeding and core hardness exceedingprior art compositions achieved at a lower cost. Variants of thedisclosed alloys are set forth in Table 3 (Nominal Compositions in wt.%):

TABLE 3 Alloy Fe Co Ni Cr Mo V C (core) A1 (modified) Bal 18 9.5 3.5 1.10.08 0.16 C3 (modified) Bal 27.8-28.2 2.9-3.1 5.0-5.2 2.4-2.6 0.18-.0250.07 CS1 (modified) Bal 15 1.5 9.0 0.2  0.08

The variants are considered equivalent to the unmodified alloys of Table1 and are within the class of alloys comprising the alloys of theinvention.

It has also been discovered that the disclosed alloys may be subjectedto a nitriding process as well as a carburizing process and other casehardening processes such as induction heating to enhance surfacehardness. For example the alloy C3 set forth above may be nitridedduring the tempering step to increase surface hardness by 10% or more.

FIGS. 10 and 11 depict the incorporation of steels of the type discussedabove in a die. Various types of dies may be fabricating using thedescribed steels, including roll form dies as well as punch dies. Sheetforming dies or stamping dies may also be made using the describedsteels.

Referring to FIG. 10, there is illustrated a representation of a rollform die 48. Draw dies used to form wire products may also be made usingthe described steels. Die 48 typically includes a shaped form or surface50 which is engaged against a material (not shown) which is to be formedor shaped by the die 48. Typically, the die 48 is incorporated in ordefines a roll and rotates as depicted by the arrow in FIG. 10 againstthe material to shape and form the material. The surface 50 desirablyhas a hardness greater than 58 Rockwell C and a case depth of preferablymore than about 0.1 millimeter. More preferably, the surface hardnessexceeds 65 Rockwell C and the depth of the case of the hardened die isgreater than about 1 millimeter and preferably in the range of about 1to 3 millimeters. The surface may be hardened by carburizing, nitriding,induction heating or other methods to improve the strength of thematerial at the surface and thereby achieve case hardening.

Importantly, the core of the die has a lower hardness in the range ofless than about 53 Rockwell C. This improves the wear and fatigueresistance at the surface while maintaining a lower hardness core whichis ductile. Thus, there is a gradient in hardness with respect to thesurface and core with the high hardness surface providing high wear andfatigue resistance while the ductile core provides shock and flawtolerance. Also, the compressive residual stresses in the case improvethe contact and bending fatigue resistance at the surface which canotherwise lead to premature failure of the tool.

FIG. 11 illustrates in cross section, a typical construction of such adie 48 wherein the surface 50 comprises the hardened portion of the die48 inasmuch as the curved or shaped surface 50 engages the materialwhich is being shaped by the die. Besides carburizing, nitriding,induction hardening, other surface modification techniques may be usedfollowed by heat treatment to produce the desired case and coreproperties. The result is a very high hardness for the surface,particularly in high wear areas as well as beneficial residualcompressive stresses. The strength of the core can also be controlled toprovide sufficient strength to withstand the internal body stressesapplied to the tool. The lower hardness of the core also providesimproved flaw tolerance and ductility. As a result, the tool hassignificantly longer useful life.

Of course, the drawing depicts a roll form die but other types of diesincluding punch dies, sheet fabrication dies, draw dies and the like maybe fabricated. It is noted that with roll form dies, the resultant dieand the process of using the die provides improved wear, notch bending,fatigue resistance and contact fatigue resistance as well as shock andflaw tolerance. Similar benefits are observed with sheet fabrication andpunch dies. The dies exhibit a longer useful life and can be used toproduce more complex shapes as well as punch or form dies with higheraspect ratio features. These benefits provide that a single die may beutilized to provide more complex shapes and cuts. The manufacture of thedies of the invention are easier because the machining of the shape ofthe tool can be completed in the soft state and then case hardened tohigh hardness, considering the distortion involved in such a process.

Various aspects of the invention may therefore be altered withoutchanging the form and scope of the invention. Thus, the invention is tobe limited only by the following claims and equivalents thereof.

What is claimed is:
 1. a tool selected from the group consisting of dies, punches, stamping dies and draw dies for forming metal alloy components comprising, in combination: a die form including a core and a surface form comprised of a steel alloy comprising about 0.008 to 0.24 percent by weight carbon, about 15 to 28 percent by weight cobalt, about 1.5 to 9.8 percent by weight nickel, and one or more additives selected from the group consisting of chromium, molybdenum and vanadium; and the balance iron having a core hardness less than about 53 Rockwell C and a surface hardness of at least a portion of the surface greater than about 58 Rockwell C of the same steel alloy as the core said surface conditioned by a process taken from the group consisting of carburizing, nitriding, induction hardening and combinations thereof to a surface depth more than about 0.1 millimeter.
 2. The tool of claim 1 wherein the hardness surface depth is about 1 to 3 millimeters.
 3. The tool of claim 1 wherein the surface hardness is at least about 65 Rockwell C.
 4. The tool of claim 1 wherein the surface hardness is in the range of about 58 to 62 Rockwell C.
 5. A roll form die for use in the manufacture of metal alloy components by means of a roll forming process, comprising, in combination: a form die including a case and a core comprised of a steel alloy material comprising about 0.008 to 0.24 percent by weight carbon, about 15 to 28 percent by weight cobalt, about 1.5 to 9.5 percent by weight nickel, and one or more additives selected from the group consisting of chromium, molybdenum and vanadium; and the balance iron, at least a portion of said case having a hardness of greater than about 58 Rockwell C effected by a case hardening process, said core having a hardness less than about 53 Rockwell C.
 6. The roll form die of claim 5 wherein said steel alloy case hardened by a process selected from the group consisting of carburizing, nitriding, induction hardening and combinations thereof.
 7. The roll pin die of claim 5 wherein the die includes a forming surface and the hardened portion of the case comprises the forming surface.
 8. The roll form die of claim 5 wherein the hardened portion of the surface comprises a component forming surface. 